Multifunctional graphene-silicone elastomer nanocomposite, method of making the same, and uses thereof

ABSTRACT

A nanocomposite composition having a silicone elastomer matrix having therein a filler loading of greater than 0.05 wt %, based on total nanocomposite weight, wherein the filler is functional graphene sheets (FGS) having a surface area of from 300 m 2 /g to 2630 m 2 /g; and a method for producing the nanocomposite and uses thereof.

REFERENCE TO RELATED APPLICATIONS

The present application is a Continuation of U.S. application Ser. No.12/945,043, filed Nov. 12, 2010, now allowed, and also claims priorityon U.S. Provisional Application Ser. No. 61/260,538, filed Nov. 12,2009, the entire contents of each of which are hereby incorporated byreference.

BACKGROUND OF THE INVENTION

Field of the Invention

The present invention relates to nanocomposites having a matrix ofsilicone elastomer with multifunctional graphene sheets as filler,methods of making the same and their use.

Description of the Related Art

The effect of filler dispersion on the mechanical properties of theresulting composite has been studied for decades but a consensus is yetto be reached. Many have suggested that maximizing filler dispersion iscrucial in achieving good mechanical properties. For example, for carbonnanotubes (CNT), Ajayan et al. suggested that load transfer can belimited when the nanotubes are slipping within the bundles.¹ The bundlesneed to be broken into individual dispersed tube segments to obtaineffective modulus increase and strengthening. Schandler et al. haveproposed that infiltrating the polymer into the interstices of thenanotube bundles can create effective load transferring and thereforemechanical reinforcement.² Similarly for inorganic fillers, Lebaron etal. have suggested that the complete dispersion of clay optimized thenumber of reinforcing elements for carrying an applied load anddeflecting cracks, allowing for tensile property improvements.³ ¹Ajayan, P. M.; Schadler, L. S.; Giannaris, C.; Rubio, A. AdvancedMaterials 2000, 12, (10), 750-² Schadler, L. S. Giannaris, S. C.;Ajayan, P. M. Applied Physics Letters 1998, 73, (26), 3842-3844³LeBaron, P. C.; Wang, Z.; Pinnavaia, T. J. Applied Clay Science 1999,15, (1-2), 11-29

Large clusters of particles can act as flaws to initiate prematuretermination of stretching.⁴ On the other hand, it has long beensuggested in the automotive tire industry that aggregated fillers aremore effective than primary particles in enhancing the modulus andtensile strength of the elastomer.⁵ At large strains, the deformationand irreversible breakdown of aggregates absorb energy, allowing thecomposite to tolerate higher amounts of stress. However, a rigorousunderstanding of the effect of breaking up initial filler agglomerateson the mechanical properties that incorporates the two aforementionedcontrasting views, is lacking. ⁴ Wilbrink, M. W. L.; Argon, A. S.;Cohen, R. E.; Weinberg, M. Polymer 2001, 42, (26), 10155-10180⁵Poovarodom, S.; Hosseinpour, D.; Berg, J. C. Industrial & EngineeringChemistry Research 2008, 47, (8), 2623-2629

In achieving the maximum effect with the minimum filler loading, it isimportant to understand the correlation between the spatial distributionof dispersed fillers and the macroscopic mechanical properties of thecomposite.^(6,7) Some understanding of the structure-propertyrelationship has been developed previously by others. A largeragglomeration of silica renders a better improvement in the Young'smodulus of the matrix.⁸ It has been shown by Akcora et al. thatself-assembled nanoparticle sheet yielded a solid-like rheologicalbehavior in polystyrene whereas well-dispersed short particle stringsdid not.⁹ However, the effect of filler assembly on the tensileproperties of the composites is not yet well-understood. ⁶ Vaia, R. A.;Maguire, J. F. Chemistry of Materials 2007, 19, (11), 2736-2751⁷ Balazs,A. C.; Emrick, T.; Russell, T. P. Science 2006, 314, (5802), 1107-1110⁸Oberdisse, J. Soft Matter 2006, 2, (1), 29-36⁹ Akcora, P.; Liu, H.;Kumar, S. K.; Moll, J.; Li, Y.; Benicewicz, B. C.; Schadler, L. S.;Acehan, D.; Panagiotopoulos, A. Z.; Pryamitsyn, V.; Ganesan, V.;Ilavsky, J.; Thiyagarajan, P.; Colby, R. H.; Douglas, J. F. NatureMaterials 2009, 8, (4), 354-U121

Another fundamental issue that has drawn much attention is the origin ofthe reinforcements of tensile properties in composites. Simultaneousimprovements in modulus, strength and elongation at break with theincorporation of fillers have been observed inpoly(methylmethacrylate),¹⁰ epoxy,¹¹ styrene-butadiene rubber,¹²polyimide,¹³ and silicone rubber.^(14,15,16,17,18) While the modulus andstrength increase with the filler concentration, the elongation at breakin some cases increases initially and then decreases above a criticalfiller concentration.^(11,13, 16,17) ¹⁰ Sui, X. M.; Wagner, H. D. NanoLetters 2009, 9, (4), 1423-1426¹¹ Tseng, C. H.; Wang, C. C.; Chen, C. Y.Chemistry of Materials 2007, 19, (2), 308-315¹² Bokobza, L.; Rahmani,M.; Belin, C.; Bruneel, J. L.; El Bounia, N. E. Journal Of PolymerScience Part B-Polymer Physics 2008, 46, (18), 1939-1951¹³ An, L.; Pan,Y. Z.; Shen, X. W.; Lu, H. B.; Yang, Y. L. Journal of MaterialsChemistry 2008, 18, (41), 4928-4941¹⁴ Aranguren, M. I.; Mora, E.;Macosko, C. W.; Saam, J. Rubber Chemistry And Technology 1994, 67, (5),820-833¹⁵ Yuan, Q. W.; Mark, J. E. Macromolecular Chemistry And Physics1999, 200, (1), 206-220¹⁶ Osman, M. A.; Atallah, A.; Muller, M.; Suter,U. W. Polymer 2001, 42, (15), 6545-6556¹⁷ Bokobza, L.; Rahmani, M.Kgk-Kautschuk Gummi Kunststoffe 2009, 62, (3), 112-117¹⁸ LeBaron, P. C.;Pinnavaia, T. J. Chemistry Of Materials 2001, 13, (10), 3760-3765

The increase in modulus is attributed to load transferring to thestiffer filler material. ^(19,20) Some understanding has been achievedin the tensile strength and elongation at break increase. Sui et al.demonstrated using transmission electron microscopy (TEM) the mechanismresponsible for the significant elongation at break increase inelectrospun CNT-poly(methyl methacrylate) (PMMA) fibers.¹⁰ In pure PMMAfiber, sparse and unstable necking was observed along the fiber undertension, followed by failure of the fiber. When 1.5 wt. % single wallcarbon nanotubes (SWCNT) were added, multiple necking was initiated butarrested by SWCNT ropes. Further stretching led to bridging by SWCNTropes, which caused a dilation effect in the fiber and an increase inthe elongation at break. The inelastic strain and energy dissipationintroduced by the necking and bridging was proposed to explain thetensile strength increase of the nanocomposite. Only one CNTconcentration was used. In the same study, millimeter-sized pure and CNTfilled PMMA films were studied and improvement in the elongation atbreak was also observed, although to a lesser extent compared to theelectrospun fibers. The improvement in the films was not addressed inthe study. ¹⁹ Hashin, Z.; Shtrikman, S. Journal Of The Mechanics AndPhysics Of Solids 1963, 11, (2), 127-140²⁰ Nielsen, L. E. Journal OfApplied Physics 1970, 41, (11), 4626-&

Load transferring to CNT has been proposed to explain the strength andelongation at break increase in epoxy.¹¹ When an amphiphilic blockcopolymer was incorporated into epoxy, elongation at break increase wasobserved.²¹ The underlying mechanisms were investigated with opticalmicroscopy and TEM. It was found that a 15 nm size spherical blockcopolymer micelle could cavitate to induce matrix shear banding. It wassuggested that the dilation effect and shear banding introduced by thecavitation led to the observed increase in the elongation at break. ²¹Liu, J.; Sue, H. J.; Thompson, Z. J.; Bates, F. S.; Dettloff, M.; Jacob,G.; Verghese, N.; Pham, H. Macromolecules 2008, 41, (20), 7616-7624

When rod-like attapulgite was incorporated into polyimide, simultaneousimprovements in modulus, strength and elongation at break wereobserved.¹³ The enhancement of the interfacial stress transfer and theresistance to crack propagation induced by attapulgite was proposed toexplain the mechanical reinforcement.

Filler agglomerates acting as defects have been proposed to explain thereversal in the elongation at break.^(11,22) Incorporation of freevolume with the filler has also been suggested to be causing thereversal effect.¹³ The addition of filler increased the free volume ordefects in nanocomposites and the resistance to crack propagation duringdeformation. Below the critical concentration, the latter effectdominated and elongation at break increased. Above the threshold, theincrease in the number of defects dominated and the elongation at breakstarted to decrease. The reversal effect was also observed with theincorporation of polystyrene-modified cadmium selenide nanoparticles topolystyrene (PS).²³ It was proposed that two competing effects determinethe elongation at break of the composite. Nanoparticles entrapped withinthe mature craze during craze widening disrupt the formation ofcross-tie fibrils by increasing the mobility of polymer segments at thecraze-bulk interface. Less cross-tie fibrils reduced the prematurerupture of the craze fibrils and increased the failure strain. On theother hand, entrapped nanoparticles also reduced the extensibility ofthe craze fibrils or the dilation effect of the craze. So the twocompeting effects led to a maximum in elongation at break of thecomposite as a function of nanoparticle concentrations. ²² Gorga, R. E.;Cohen, R. E. Journal of Polymer Science Part B-Polymer Physics 2004, 42,(14), 2690-2702²³ Lee, J. Y.; Zhang, Q. L.; Wang, J. Y.; Emrick, T.;Crosby, A. J. Macromolecules 2007, 40, (17), 6406-6412

The simultaneous improvements are not limited to polymeric matrices. Theincorporation of polymeric fibers increased the strength and elongationat break of the newly engineered building material called engineeredcementitious composites (ECC).²⁴ ECCs have been designed to distributemany cracks of small width throughout the composite rather than only afew large cracks seen in traditional concrete failure. Such adistributed deformation is responsible for the observed mechanicalreinforcement. Similar mechanisms have been shown to cause theelongation at break increase in biological composites such as nacre.²⁵²⁴ Li, V. C.; Wang, 5. X.; Wu, C. Aci Materials Journal 2001, 98, (6),483-492²⁵ Wang, R. Z.; Suo, Z.; Evans, A. G.; Yao, N.; Aksay, I. A.Journal Of Materials Research 2001, 16, (9), 2485-2493

Despite the aforementioned efforts, some fundamental issues governingthe tensile properties improvements have not been completely understood.For example, it is not known how the filler agglomeration and fillerconcentration influence the interaction between fillers and tears orcracks, nor how filler length scale influences the interaction. Further,it is not known how the interaction is related to the reversal effect orhow the local deformation is directly correlated with the macroscopictensile properties in bulk composites. Lastly, it is not known howmechanical load is being transferred to the filler. These are allcritical questions that need to be addressed in order to gain a completeunderstanding of the reinforcement.

One potential filler that has been suggested is functional graphenesheets (FGS). FGS is an atomically thin layer of graphite hundreds ofnanometers in the lateral dimension and decorated with carboxyls at theedges and hydroxyls and epoxides on the planes. Our group invented amethod to produce functionalized graphene sheet (FGS) on a large scale;see U.S. Patent Application Publication 2007/0092432, filed Oct. 14,2005 and published Apr. 26, 2007 (the entire contents of which arehereby incorporated by reference; hereafter “the '432 application”). Ithas a wrinkled geometry with an average aspect ratio of 500 and asurface area from 300 m²/g to 2630 m²/g, typically up to 1800m²/g.^(26,27) It is preferably produced through thermal exfoliation andreduction of oxidized natural graphite. The '432 application furtherdiscloses these FGS products. Stankovich et al. developed an alternativemethod to produce graphene.²⁸ Graphene oxide was first obtained byoxidation of natural graphite and sonication of graphite oxide. Chemicalreduction of graphene oxide yielded graphene with good electricalconductivity. In a recent study, significant increases in glasstransition temperature. Young's modulus, tensile strength and electricalconductivity was observed in when 1 weight % of FGS was incorporatedinto poly(methyl methacrylate) and poly(acrylonitrile).²⁹ An enhancementin the modulus and electrical conductivity as well as a reduction in thecoefficient of thermal expansion and gas permeability was observed whenFGS was added to poly(ethylene-2,6-naphthalate) andpoly(carbonate).^(30,31) When reduced graphene oxide was incorporatedinto polystyrene, a low electrical percolation of 0.1 vol. % and goodconductivities were obtained.²⁸ ²⁶ Schniepp, H. C.; Kudin, K. N.; Li, J.L.; Prud′homme, R. K.; Car, R.; Saville, D. A.; Aksay, I. A. Acs Nano2008, 2, (12), 2577-2584²⁷ McAllister, M. J.; Li, J. L.; Adamson, D. H.;Schniepp, H. C.; Abdala, A. A.; Liu, J.; Herrera-Alonso, M.; Milius, D.L.; Car, R.; Prud'homme, R. K.; Aksay, I. A. Chemistry Of Materials2007, 19, (18), 4396-4404²⁸ Stankovich, S.; Dikin, D. A.; Dommett, G. H.B.; Kohlhaas, K. M.; Zimney, E. J.; Stach, E. A.; Piner, R. D.; Nguyen,S. T.; Ruoff, R. S. Nature 2006, 442, (7100), 282-286²⁹ Ramanathan, T.;Abdala, A. A.; Stankovich, S.; Dikin, D. A.; Herrera-Alonso, M.; Piner,R. D.; Adamson, D. H.; Schniepp, H. C.; Chen, X.; Ruoff, R. S.; Nguyen,S. T.; Aksay, I. A.; Prud'homme, R. K.; Brinson, L. C. NatureNanotechnology 2008, 3, (6), 327-331³⁰ Kim, H.; Macosko, C. W.Macromolecules 2008, 41, (9), 3317-3327³¹ Kim, H.; Macosko, C. W.Polymer 2009, 50, (15), 3797-3809

U.S. patent application Ser. No. 11/543,872, filed Oct. 6, 2006 (theentire contents of which are hereby incorporated by reference),discloses the use of the FGS of the '432 application in the productionof various nanocomposite rubbers.

SE has attracted both scientific and commercial interest for its thermalstability over a wide range of temperatures (−50 to over 200° C.),retention of elastomeric properties at low temperatures due to a lowglass transition temperature of −125° C., its chemical and weatheringresistance.^(32,33,34) SE is typically made by end-linking poly(dimethylsiloxane) (PDMS) and therefore its molecular weight between crosslinksis well-characterized. Due to its relatively inferior tensile strengthin the unfilled state (typically less than 1 MPa, compared to more than10 MPa of natural rubber), silica is generally used to render SEapplicable in commercial applications.^(34,34) Other fillers includingsilica,^(14,15,35) s clays,^(16,16,36) carbon nanotubes (CNT),^(17,37)graphite nanosheet,³⁸ glass fiber,³⁹ and in-situ precipitated alumina,⁴⁰have also been studied as alternative fillers for SE. ³² Mark, J. E.Accounts Of Chemical Research 2004, 37, (12), 946-953³³ Noll, W.,Chemistry and Technology of Silicones. Academic Press, Inc.: New York,1978³⁴ Butts, M.; et. al. In Kirk-Othmer Encyclopedia of ChemicalTechnology-Silicones. Wiley Interscience: New York, 2004³⁵ Mark, J. E.;Jiang, C. Y.; Tang, M. Y. Macromolecules 1984, 17, (12), 2613-2616³⁶Osman, M. A.; Atallah, A.; Kahr, G.; Suter, U. W. Journal of AppliedPolymer Science 2002, 83, (10), 2175-2183³⁷ Frogley, M. D.; Ravich, D.;Wagner, H. D. Composites Science And Technology 2003, 63, (11),1647-1654³⁸ Chen, L.; Lu, L.; Wu, D. J.; Chen, G. H. Polymer Composites2007, 28, (4), 493-498³⁹ Park, E. S. Journal of Applied Polymer Science2007, 105, (2), 460-468⁴⁰ Mark, J. E.; Wang, S. B. Polymer Bulletin1988, 20, (5), 443-448

SUMMARY OF THE INVENTION

Accordingly, one object of the present invention is to provide ananocomposite based on silicone elastomers that has one or more ofhigher modulus, strength, failure strain, electrical conductivity andlower gas permeability than the unfilled silicone elastomer.

A further object of the present invention is to provide a method forproducing such a nanocomposite.

A further object of the present invention is to provide articles madefrom the nanocomposite, including, but not limited to electricallyconductive and low-permeability coating, adhesive and sealants, as wellas flexible electrodes, actuators, pressure sensor, printed circuits andelectromagnetic interference shielding material.

These and other objects of the present invention, either alone or incombinations thereof, have been satisfied by the discovery of ananocomposite composition comprising:

a silicone elastomer matrix having therein a filler loading of greaterthan 0.05 wt %, based on total nanocomposite weight;

wherein the filler is functional graphene sheets (FGS) having a surfacearea of from 300 m²/g to 2630 m²/g;

a method for producting the nanocomposite composition and its use in avariety of end products.

BRIEF DESCRIPTION OF THE DRAWINGS

The patent or application file contains at least one drawing executed incolor. Copies of this patent or patent application publication withcolor drawing(s) will be provided by the Office upon request and paymentof the necessary fee.

A more complete appreciation of the invention and many of the attendantadvantages thereof will be readily obtained as the same becomes betterunderstood by reference to the following detailed description whenconsidered in connection with the accompanying drawings, wherein:

FIGS. 1A-IF show SEM images of cryo-fractured unfilled and FGS-filled SEsurfaces.

FIG. 2 provides a graphical representation of the effect of fillerconcentration on the electrical conductivity of various FGS and graphenefilled nanocomposites.

FIG. 3 provides a graphical representation of stress-strain curves ofFGS-SE nanocomposites at different FGS concentrations.

FIGS. 4A-4F provide photographs of the tearing of unfilled SE (A) and0.5 wt. % FGS-SE (B)-(D) and SEM images of tensile-fractured surfaces ofunfilled (E) and 0.5 wt % FGS-filled SE (F).

FIG. 5 provides a graphical representation of hysteresischaracterization of unfilled and 0.5 wt % FGS-SE nanocomposite.

FIGS. 6A-6D provide images of the deformed lattice in unfilled andFGS-SE nanocomposites.

FIGS. 7A-7C provide graphical representations of (7A). Simulatednormalized stress-strain curves of unfilled and FGS-filled SE; (7B).Fraction of the matrix torn versus FGS vol. % at three differentstrains; and (7C). Average strain of tears versus FGS vol. %.

FIGS. 8A and 8B provide graphical representations of a comparison of FGSwith other fillers in the modulus of the composite and the improvementin the modulus.^(14,16,17,18,36,38)

FIGS. 9A and 9B provide graphical representations of a comparison of FGSwith other fillers in the tensile strength of the composite and theimprovement in tensile strength rendered by the filler.^(14-18,36,39,40)

FIG. 10 provides a graphical representation of the effect of catalystconcentration on the modulus of SE at r=1.5 for all samples.

FIG. 11 provides a graphical representation of the effect of siliconhydride to vinyl ratio on the modulus of FGS-SE nanocomposite.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

The present invention involves the addition of functionalized graphenesheets (FGS) to silicone elastomer. Within the context of the presentinvention the term “silicone elastomer” is used to refer to any of avariety of elastomeric modified silicone polymers as distinct fromunmodified polydimethylsiloxane (PDMS). The invention has highermodulus, strength, failure strain, electrical conductivity and lower gaspermeability than the unfilled silicone elastomer. The current inventionintroduces new applications for silicone elastomer, such as electricallyconductive and low-permeability coating, adhesive and sealants, as wellas flexible electrodes, actuators, pressure sensor, printed circuits andelectromagnetic interference shielding material.

The FGS-silicone elastomer nanocomposite of a most preferred embodimentof the present invention simultaneously has superior mechanical,electrical and barrier properties compared to unfilled siliconeelastomers. Furthermore, the current invention provides a method todetect internal damage within the material through conductancemeasurements. So when used in the industry, it possesseshealth-monitoring capability which can be crucial for applications usingthe nanocomposite. The product also has lower density than commerciallyavailable silicone and therefore can reduce the energy cost associatedwith transporting and using the product.

FGS can be produced via a process that has been described in publishedarticles (H. C. Schniepp, J.-L. Li, M. J. McAllister, et al., J. Phys.Chem. B 110, 8535-39, 2006; M. J. McAllister, J.-L. Li, D. H. Adamson,H. C. Schniepp, et al., Chem. Materials 19, 4396-4404, 2007) (the entirecontents of each of which are hereby incorporated by reference) and the'432 application noted above.

In a preferred embodiment of the present invention method for formingthe nanocomposite, FGS is dispersed in a polar solvent, such astetrahydrofuran, and probe-sonicated. Then the suspension is combinedwith a vinyl terminated polysiloxane (preferably a vinyl terminatedpolydimethylsiloxane) and the polar solvent is completely evaporatedoff. An appropriate crosslinker and hydrosilylation catalyst (preferablya platinum complex catalyst) are combined with the resulting mixture andthe mixture is cured at elevated temperature, preferably about 100° C.for a period of time from 1 to 48 hours, preferably from 5 to 30 hours,more preferably from 20-25 hours, most preferably approximately 24hours.

The crosslinking reaction for the silicone elastomer involves thereaction between the crosslinker, and the vinyl terminating groups onthe vinyl-terminated polysiloxane in the presence of a hydrosilylationcatalyst. Suitable crosslinking agents include any conventionalcrosslinking agent, such as those disclosed in “The Basics of SiliconChemistry” (Dow Corning Publication); W. Noll, Chemistry and technologyof Silicones, Academic Press, New York (1968); T. C. Kendrick, B.Parbhoo, J. W. White, “Siloxane Polymers and Copolymers,” in TheChemistry of Organic Silicon Compounds Pt. 2 (edited by S. Patai and Z.Rappoport), 21, p. 1289-1361, John Wiley, Chichester (1989); and S. J.Clarson, J. A. Semlyen, Siloxane Polymers, Prentice Hall, New Jersey(1993), the contents of each of which are hereby incorporated byreference. Preferably, the crosslinking agent is selected from tetrakis(dimethyl siloxy) silanes, or poly(hydromethylsiloxane) crosslinkers.The resulting mechanical properties, electrical properties and gaspermeability of the FGS-silicone elastomer nanocomposite showedincreased modulus, elongation at break, tensile strength and electricalconductivity and decreased gas permeability, as compared to the samesilicone elastomer without the FGS filler.

The hydrosilylation catalyst is not particularly restricted, and can beany conventional hydrosilylation catalyst. Specific examples include,but are not limited to, chloroplatinic acid, elementary platinum, solidplatinum supported on a carrier such as alumina, silica or carbon black;platinum-vinylsiloxane complexes {e.g. Pt_(n) (ViMe₂SiOSiMe₂Vi)_(n),Pt[(MeViSiO)₄]_(m)}; platinum-phosphine complexes {e.g. Pt(PPh₃)₄,Pt(PBU₃)₄}; platinum-phosphite complexes {e.g. Pt[P(OPh)₃]₄,Pt[P(OBu)₃]₄} (in the above formulas, Me stands for methyl, Bu forbutyl, Vi for vinyl, Ph for phenyl, and n and m each represents aninteger); Pt (acac)₂; and platinum-hydrocarbon conjugates described byAshby et al. in U.S. Pat. Nos. 3,159,601 and 3,159,662 as well asplatinum alcoholates described by Lamoreaux et al. in U.S. Pat. No.3,220,972, the contents of each of which are hereby incorporated byreference.

As examples of the catalyst, other than platinum compounds, there may bementioned RhCl(PPh₃)₃, RhCl₃, Rh/Al₂O₃, RuCl₃, IrCl₃, FeCl₃, AlCl₃,PdCl₂.2H₂O, NiCl₂, TiCl₄, etc. These catalysts may be used singly or twoor more of them may be used in combination. From the viewpoint ofcatalytic activity, chloroplatinic acid, platinum-olefin complexes,platinum-vinylsiloxane complexes, Pt(acac), and the like are preferred,with platinum-cyclovinylmethylsiloxane complex being most preferred. Theamount of the catalyst is not particularly restricted but the catalystis preferably used in an amount within the range of 10⁻¹ to 10⁻⁸ moles,more preferably 10⁻² to 10⁻⁶ moles, per mole of the alkenyl group in thevinyl-terminated polysiloxane. Hydrosilylation catalysts are generallyexpensive and corrosive and, in some instances, they induce generationof hydrogen gas in large amount to thereby cause foaming of curedproducts. Therefore, it is recommended that their use in an amount ofmore than 10⁻¹ moles be avoided.

Within the context of the present invention, the term “vinyl-terminatedpolysiloxane” is used to represent a component of the present inventionsiloxane elastomer that contains at least one diorganosiloxane unit andhas at least two silicon-bonded alkenyl groups in each molecule. Thealkenyl group can be exemplified by vinyl, allyl, butenyl, pentenyl,hexenyl, and heptenyl and is preferably vinyl. The non-alkenyl Si-bondedorganic groups are exemplified by alkyl groups such as methyl, ethyl,propyl, butyl, pentyl, and hexyl; aryl groups such as phenyl, tolyl, andxylyl; and halogenated alkyl groups such as 3-chloropropyl and3,3,3-trifluoropropyl, and is preferably methyl and/or phenyl. Themolecular structure of the vinyl-terminated polysiloxane is not criticalas long as it contains at least one diorganosiloxane unit, i.e.,siloxane unit with a general formula R₂SiO_(2/2). As other siloxaneunits, the vinyl-terminated polysiloxane may contain small amounts ofsiloxane unit with a general formula R₃SiO_(1/2), siloxane unit with ageneral formula RSiO_(3/2), and siloxane unit with a general formula.SiO_(4/2). R in the preceding formulas represents a substituted orunsubstituted monovalent hydrocarbon group and can be exemplified by thealkyl, alkenyl, aryl, and halogenated alkyl referenced above. Themolecular structure of the vinyl-terminated polysiloxane can beexemplified by straight chain, branched chain, partially branchedstraight chain, and dendritic, wherein straight chain, branched chain,and partially branched straight chain are preferred. The viscosity ofthe vinyl-terminated polysiloxane at 25° C. is not critical, but ispreferably 100 to 1,000,000 mPas and more preferably is 100 to 500,000mPas, most preferably from 100 to 300,000 mPas. The weight averagemolecular weight of the vinyl-terminated polysiloxane is also notparticularly critical and will depend on the end use desired for thefinished FGS-SE composition. Preferably the weight average molecularweight of the vinyl-terminated polysiloxane is in a range from 5000 to2,000,000, more preferably from 5000 to 50,000, most preferably from8000 to 12,000.

The vinyl-terminated polysiloxane is preferably a member selected fromdimethylvinylsiloxy-endblocked dimethylpolysiloxanes;dimethylvinylsiloxy-endblocked dimethylsiloxane-methylvinylsiloxanecopolymers; trimethylsiloxy-endblockeddimethylsiloxane-methylvinylsiloxane copolymers; branched-chaindimethylpolysiloxane with molecular chain ends terminated bydimethylvinylsiloxy and trimethylsiloxy; trimethylsiloxy-endblockedbranched-chain dimethylsiloxane-methylvinylsiloxane copolymers; theorganopolysiloxanes afforded by replacing all or part of the methyl inthe preceding organopolysiloxanes with alkyl such as ethyl or propyl,aryl such as phenyl or tolyl, or halogenated alkyl such as3,3,3-trifluoropropyl; the organopolysiloxanes afforded by replacing allor part of the vinyl in the preceding organopolysiloxanes with alkenylsuch as allyl or propenyl; and mixtures of two or more of the precedingorganopolysiloxanes.

For convenience, the vinyl-terminated polysiloxane will be discussedwith reference to a vinyl-terminated poly(dimethylsiloxane). However,this is not intended to be limiting of the present invention, but merelyused in an exemplary manner for convenience.

The present invention nanocomposite properties provide the ability tomonitor the structural health of products formed from the nanocompositeby measuring conductance properties to detect internal damage in theresulting product.

In the product of the present invention, tensile properties improvementsare preferably achieved when FGS is percolated in SE. Within the contextof the present invention, the term “percolated” is intended to indicatethat a continuous path is established in three dimensions through theFGS by formation of a connected FGS network with nanometer scaleseparation at the contact point between individual sheets. Normally, theFGS sheets are statistically in contact. Indications of percolation arethe onset of the transition from non-electrically conducting toelectrically conducting, or the state in which the storage and lossmoduli measured as a function of frequency (G′(ω) and G′(ω),respectively) scale as G′(ω)˜G′(ω)˜ω^(n). These characteristics aremeant to be indicative of percolation and are not intended as limitingthe present invention. Agglomeration of FGS can be observed using SEM.The agglomeration facilitates electrical percolation and thereforetensile properties improvements. Although the present inventors do notwish to be bound by any particular mechanistic explanation for theimprovement in properties in the present invention, it is believed thatthe increase of tensile strength can be attributed to load transfer toFGS. The increase of elongation at break is believed to be due to thedilation effect of tearing and distributed deformation introduced by thepercolated FGS network. The reversal in the elongation at break isobserved and is believed to be due to the competing effects of thedegree of tear opening and the number of tears with increasing FGSconcentration. Multifunctional reinforcement of SE by FGS is alsodemonstrated.

Experimental Section

2.1. Materials.

Vinyl-terminated PDMS with an average molecular weight of 9400,tetrakis(dimethylsiloxy)silane and platinum-cyclovinylmethylsiloxanecomplex were obtained from Gelest, Inc. Tetrahydrofuran (THF) waspurchased from Sigma Aldrich. FGS was produced using a thermalexfoliation method previously reported using graphite oxide (GO)supplied by Vorbeck Materials.^(27,27) The carbon to oxygen ratio of theFGS was determined to be 15 to 1 using modified classical Pregl andDumas method by Atlantic Microlab, Inc.⁴¹ ⁴¹ Patterson, R. K. AnalyticalChemistry 1973, 45, (3), 605-609

2.2. Processing of Unfilled SE and FGS-SE Nanocomposite.

An SE network was prepared by end-linking the di-functionalvinyl-terminated PDMS molecules and the tetra-functional crosslinkertetrakis(dimethylsiloxy silane) with platinum cyclovinylmethylsiloxanecomplex as the catalyst. The crosslinking resulted from the reaction ofterminating vinyl groups on the PDMS with silicon hydride groups on thetetrakis(dimethylsiloxy silane).

The unfilled SE samples were produced as follows: predetermined amountsof PDMS, the crosslinking agent tetrakis(dimethylsiloxy)silane and thecatalyst platinum-cyclovinylmethylsiloxane were mixed by magneticstirring for 20 min; the mixture was then poured onto apolytetrafluoroethylene mold and cured at 100° C. for 12 h. FGS-SEnanocomposites were produced as follows: an FGS suspension with aconcentration of 1 mg/ml was made by mixing a predetermined amount ofFGS and tetrahydrofuran (THF) in a beaker. The beaker was immersed in anice bath while the suspension was probe-sonicated for 30 mum (VirSonic100, The Virtis Co., NY; with an output power 12 W). After sonication,the suspension was transferred to another beaker containing a desiredamount of PDMS polymer. The mixture containing the FGS, THF, and PDMSwas placed on a stir plate heated to 60° C. to evaporate off all the THFwith magnetic stirring. After all the THF evaporated, the thixotropicmixture was cooled to room temperature beforetetrakis(dimethylsiloxy)silane and platinum-cyclovinylmethylsiloxanewere added. The mixture was hand-mixed with a steel spatula for 15 min.The final mixture was then transferred to a polytetrafluoroethylenemold. A metal plate was used to shear and spread the mixture evenlyacross the mold. The shearing velocity of the plate was 6 cm/s. Finally,the mixture was cured in an oven at 100° C. for 12 h. For 3 wt % FGS-SEnanocomposites, the samples were prepared using vacuum molding tominimize trapped air bubbles.

2.3. Mechanical Property Measurements.

Tensile and mechanical hysteresis measurements were made under ambientconditions using an Instron tensile testing machine (Model II22,Instron, MA). The dog-bone-shaped samples used in the measurements were22.55 mm long and 4.55 mm wide in the narrow region. Thickness of thesamples varied between 0.2 to 0.6 mm. The strain rate was set to 50.8mm/min. For the hysteresis measurements, samples were stretched to70%-80% of its average failure strain, returned to a stress level ofzero and were stretched again to a strain level similar to that of thefirst stretch. Samples were then placed in an oven set to 100° C. torecover for 24 h and then their stress-strain curves were measuredagain. The area under the stress-strain curve was calculated and thedifference in the area between the first stretch and the stretch afterrecovery was obtained. Hysteresis loss ratio was computed by dividingthe difference with the area of the first stretch. The reportedhysteresis loss ratio is an average from three samples.

2.4. Scanning Electron Microscopy Characterization (SEM).

Images of cryo-fractured SE or FGS-filled SE were taken with twodifferent SEMs. Tescan Vega SEM (Tescan USA, PA) was used tocharacterize the sample without conductive coatings at magnifications upto 3700. To obtain high resolution images, the samples were coated with3 nm iridium. An FEI XL-30 field emission gun SEM (Philips, MA) was usedto image the samples.

2.5. Electrical Conductivity Measurements.

The direct current transverse resistivity (the resistivity across thefilm thickness direction) of the FGS-SE nanocomposites was measured witha resistivity test fixture (Keithley 8009, Keithley Instrument Inc., OH)coupled with a digital multimeter (Keithley 6517). The composite filmwas cut into a circular film with a diameter of 70 mm and placed betweenthe top and guarded electrodes for the measurement. The DC longitudinalresistivity (the resistivity along the in-plane direction of the film)was measured using a standard 4-point technique. The nanocomposite filmwas cut into rectangular shape films (1-2 cm in width and 2-4 cm inlength). A film was placed on a polystyrene petri-dish and conductivecopper-nickel adhesive tape (Electron Microscopy Sciences) was placednear the two ends of the film. Conductive carbon paste (ElectronMicroscopy Science) was used to draw conductive paths between the sampleand the copper tape. The resistance was measured with a DC power supply(Tektronix PS2521G, Tektronix, OR), digital multimeter (Fluke 27, FlukeCorporation, WA) and electrometer (Keithley 6514). The conductivity of asample film was calculated based on the dimension of the film. Thelongitudinal conductivity of SE with FGS concentration less than 0.2 wt.% was below the detection limit of the devices and thus could not bemeasured. All the electrical conductivities were the average from twoseparately made samples.

2.6. Gas Permeation Measurements.

Oxygen and nitrogen permeability of unfilled and FGS-filled SE wasobtained using a constant pressure/variable volume type permeation cellfrom Professor Donald Paul's lab at University of Texas.⁴² The amount ofgas that has permeated was measured and plotted as a function of time.The permeability was determined from the slope of the linear portion ofthe plot (steady state). ⁴² Takahashi, S.; Goldberg, H. A.; Feeney, C.A.; Karim, D. P.; Farrell, M.; O'Leary, K.; Paul, D. R. Polymer 2006,47, (9), 3083-3093

2.7. Two Dimensional Viscoelastic Lattice Model.

A two dimensional viscoelastic lattice model for the elastomer matrixwith the ability to visualize tearing was utilized to explain themechanical reinforcement in FGS-SE nanocomposites. Detailed descriptionof the model is provided elsewhere.⁴³ Briefly, the model is composed ofone dimensional trusses arranged in a two dimensional (the thirddimension is of unit thickness) triangular lattice. A Zener viscoelasticelement is used to model the behavior of each truss. A tear can beinitiated when the axial stress of one truss element exceeds aprescribed breaking stress. It is known that like other materials,elastomers have intrinsic defects tens to hundreds of microns in size,that are possibly introduced while molding or cutting a testsample.^(44,45) Tearing is first initiated from these defects. Due tothe presence of these weak links, other parts of the matrix may not evenbe sampled mechanically when the failure occurs. Such heterogeneity ofbreaking stress within the SE matrix is incorporated in the model byassigning spatially varying breaking stress across the matrix. ⁴³Sanborn, S. E.; Pan, S.; Prevost, J. H.; Aksay, I. A. Submitted toMacromolecules ⁴⁴ Choi, I. S.; Roland, C. M. Rubber Chemistry AndTechnology 1996, 69, (4), 591-599⁴⁵ Hamed, G. R. Rubber Chemistry AndTechnology 1983, 56, (1), 244-251

Two domains were simulated in the model: a small length-scale model withunpercolated FGS and a large length-scale model with percolated FGSnetwork. The number of trusses was kept constant. In the smalllength-scale model, the representative volume element (RYE) had asimilar length-scale to that of the weak links in the matrix, whosedimension was set to be an order of magnitude larger than the lengthscale of individual FGS. In the large length-scale model, the percolatedFGS had a length scale comparable to the RYE. The matrix was set to havehomogeneous breaking stress since the percolated FGS has much largerlength scale than the heterogeneities. In the FGS-SE nanocompositemodel, individual FGS and percolated FGS were represented by black lineswith a stiffness four orders of magnitude larger than that of thematrix. FGS itself and the FGS-SE interface does not fail in the model.One hundred simulations were run for each FGS concentration.

The matrix is deformed in the tensile direction at a strain rate of0.0076/s, which is the loading rate used in the experiment.

Results and Discussion

3.1 Characterization of FGS Dispersion.

To elucidate the effect of filler agglomeration on the mechanicalproperties of FGS-SE nanocomposites, SEM was used to characterize theFGS dispersion state in SE matrix. The images of cryo-fractured surfacesof unfilled and FGS-filled SE are shown in FIGS. 1A-1F. Thecryo-fractured surface of the unfilled SE without conductive coating wassmooth (FIG. 1A). The morphology of FGS-filled SE was very differentfrom that of the unfilled SE. As shown in FIG. 1B, the back-scatteredelectron SEM image of an uncoated cryo-fractured-surface of 0.2 wt. %FGS-SE showed the presence of rough and smooth morphologies. The roughmorphology was likely due to the presence of FGS. In the secondaryelectron image of the same area (FIG. 1C), both bright and dark regionswere observed. The dark regions correlated well with the rough regionsin the back-scattered electron image. As the dark regions did not appearin the same sample with a conductive coating, they are due to theconductivity variation across the fractured surface. In SEM imaging, anelectron beam bombards the sample and regions with low electricalconductivity or without conductive pathways would accumulate charges dueto the lack of charge dissipation mechanism and therefore appearbrighter in the image. When regions with spatially varyingconductivities exist in a sample, regions with higher conductivity wouldappear to be darker than less-conducting regions. Since FGS was the onlyfiller in the nanocomposite, the dark regions must be the percolatedFGS-rich regions and the bright regions were the FGS-lean regions. Asshown in FIG. 1D, the wrinkled morphology of the FGS rich regionsresembled that of the FGS, confirming the agglomeration of FGSs.

The above evidence suggests that at a 0.2 wt. % FGS loading, we have acomposite material at two length scales: the first one is thesegregation of FGS-rich and FGS-lean regions with a length scale of 5-15μm and the second length scale is the ultimate FGS-SE nanocomposite inthe FGS-rich regions.

At a high enough FGS concentration, the entire sample is expected to becomposed of FGS-rich regions. That was indeed observed. When the FGSconcentration was increased to 0.8 wt. %, the conductivity inducedcontrast in the SEM image disappeared, indicating the existence ofFGS-rich regions across the entire sample (FIG. 1E). The uniformdispersion of FGS in 1 wt. % FGS-SE was demonstrated in FIG. 1F.

3.2 Electrical Properties of FGS-SE Nanocomposite.

To characterize the percolation threshold of FGS in SE, the electricalconductivity as a function of FGS weight percentage was measured and isshown in Error! Reference source not found. At 0.05 wt. % loading, therewas no increase in the transverse electrical conductivity, indicating apercolated FGS network had not yet formed. The transverse conductivityincreased by almost 6 orders of magnitude when 0.1 wt. % FGS was added.At an FGS loading of only 0.2 wt. %, the longitudinal conductivityincreased by more than 10 orders of magnitude and the transverseconductivity increased by 7 orders of magnitude. Anisotropy inelectrical conductivity was observed. The longitudinal conductivity at0.2 wt. % loading has already satisfied the conductivity requirement forelectrostatic dissipation (10⁻⁵ S/m) and also the electrostatic paintingapplications (10⁻⁴ S/m).⁴⁶ At 0.5 wt. % loading, the electricalconductivity of SE increased to 4.5×10⁻³ S/m in the longitudinaldirection and to 1.6×10⁻⁶ S/m in the transverse direction. Furtherincrease in the FGS loading above 0.5 wt. % led to a more gradualenhancement in conductivity. The conductivity reached 0.89 S/m in thelongitudinal direction and 6.6×10⁻⁴ S/m in the transverse direction at 3wt. % loading. ⁴⁶ Ramasubramaniam, R.; Chen, J.; Liu, H. Y. AppliedPhysics Letters 2003, 83, (14), 2928-2930

When conductive fillers form a network of connected paths through theinsulating matrix, a rapid increase in the electrical conductivity isexpected.⁴⁷ FIGS. 1A-1F suggest that the percolation threshold wasbetween 0.05 wt. % and 0.1 wt. % as evidenced by a rapid increase in thetranverse conductivity (almost 6 orders of magnitude) followed by a moregradual increase in conductivity (1 order of magnitude increase from 0.1wt. % to 0.2 wt. % FGS). The observed electrical percolation thresholdis, to the best of our knowledge, among the lowest in filled SE, secondonly to one case of multiwall carbon nanotube (MWNT)-SE nanocomposite.¹⁷The conductivity of FGS-SE as a function of filler concentration iscompared to that of other graphene-based polymer nanocomposites, asshown in Error! Reference source not found.A-1F. The observed electricalpercolation in FGS-SE is lower than that of graphene-polymernanocomposites previously reported.^(28,30,31,48,49,50,51) Thelongitudinal conductivities of FGS-SE are comparable to the best MWNTfilled SE and graphene based polymer nanocomposites.^(28,52) ⁴⁷ Ruschau,G. R.; Yoshikawa, S.; Newnham, R. E. Journal Of Applied Physics 1992,72, (3), 953-959⁴⁸ Ansari, S.; Gianneiis, E. P. Journal of PolymerScience Part B-Polymer Physics 2009, 47, (9), 888-897⁴⁹ Nguyen, D. A.;Lee, Y. R.; Raghu, A. V.; Jeong, H. M.; Shin, C. M.; Kim, B. K. PolymerInternational 2009, 58, (4), 412-417⁵⁰ Liang, J. J.; Wang, Y.; Huang,Y.; Ma, Y. F.; Liu, Z. F.; Cai, F. M.; Zhang, C. D.; Gao, H. J.; Chen,Y. S. Carbon 2009, 47, (3), 922-925⁵¹ Steurer, P.; Wissert, R.; Thomann,R.; Muihaupt, R. Macromolecular Rapid Communications 2009, 30, (4-5),316-327⁵² Khosla, A.; Gray, B. L. Materials Letters 2009, 63, (13-14),1203-1206

The electrical percolation threshold is influenced by the filler aspectratio and shape, as well as filler dispersion in the matrix.^(53,54,55)A theoretical percolation threshold of plates having an aspect ratio of476 is estimated to be 0.27 vol. %.⁵³ The experimentally observedpercolation was less than one fourth of the theoretically predictedvalue. The lower percolation threshold can be attributed to theagglomeration of FGS which lowers the percolation threshold.^(54,55)There exists van der Waals' attraction between FGSs. Since PDMS hasattractive interaction with FGS by forming hydrogen bonds, PDMS chainscan introduce bridging attraction. Both van der Waals's and bridgingattraction contribute to the agglomeration of FGS. In another study byour group, homogeneous dispersion of FGS in poly(ethylene oxide) led toa higher percolation threshold of 1 wt. %, corroborating theagglomeration-induced percolation of FGS in SE.⁵⁶ Due to the FGSagglomeration, the percolation threshold of FGS-SE was not determinedusing a typical percolation model which assumes homogeneous distributionof fillers.⁵⁷ ⁵³ Garboczi, E. J.; Snyder, K. A.; Douglas, J. F.; Thorpe,M. F. Physical Review E 1995, 52, (1), 819-828⁵⁴ Pegel, S.; Potschke,P.; Petzold, G.; Alig, I.; Dudkin, S. M.; Lellinger, D. Polymer 2008,49, (4), 974-984.⁵⁵ Alig, I.; Lellinger, D.; Engel, M.; Skipa, T.;Potschke, P. Polymer 2008, 49, (7), 1902-1909⁵⁶ Korkut, S.; et. al.Manuscript in preparation.⁵⁷ McLachlan, D. S.; Chiteme, C.; Park, C.;Wise, K. E.; Lowther, S. E.; Lillehei, P. T.; Siochi, E. J.; Harrison,J. S. Journal Of Polymer Science Part B-Polymer Physics 2005, 43, (22),3273-3287

The electrical conductivity of a conductive composite is governed by theintrinsic conductivity of fillers, constriction and tunneling resistanceat the contact between fillers and the number of contact spots.^(47,58)Constriction resistance is associated with constriction of electron flowthrough the contact area between two filler particles and is inverselyproportional to the contact area. The morphology of graphene sheet,which is determined by its functional groups and defects,⁵⁹ caninfluence the contact area and therefore the constriction resistance inthe composites. ⁵⁸ Simmons, J. G. Journal Of Applied Physics 1963, 34,(6), 1793-&⁵⁹ Schniepp, H. C.; Kudin, K. N.; Li, J. L.; Prud′homme, R.K.; Car, R.; Saville, D. A.; Aksay, I. A. Acs Nano 2008, 2, (12),2577-2584

Tunneling resistance is due to the tunneling of electrons throughinsulating films covering the fillers and it is proportional to the workfunction of the conductor, thickness and dielectric and thermalproperties of the film.^(47,58) The dielectric constant of the matrixinfluences the barrier height, distance of tunneling and therefore thetunneling resistance. Since the dielectric constant of most materials isa function of temperature, thermal properties of the matrix also plays arole in the tunneling resistance.⁶⁰ The thermal expansion coefficient ofthe matrix is also important. Contact force between the filler, which isdetermined by the internal stress inside the composite, stronglyinfluences the tunneling distance and therefore the overall conductivityof the composite.⁶¹ During heat curing of SE and subsequent cooling toroom temperature, volumetric shrinkage of SE could occur which inducedcompressive stress between FGS. Shrinkage from the processing ofcomposites can alter the tunneling distance.⁶² ⁶⁰ von Hippel, A. R.,Dielectric Materials and Applications. Technology Press of MIT:Cambridge, 1961⁶¹ Li, L.; Morris, J. E. In Electrical conduction modelsfor isotropically conductive adhesive joints, 1997; Ieee-Inst ElectricalElectronics Engineers Inc: 1997; pp 3-8⁶² Zweifel, Y.; Plummer, C. J.G.; Kausch, H. H. Journal Of Materials Science 1998, 33, (7), 1715-1721

The number of contact spots between fillers is influenced by theirdispersion. Better dispersion enables more contacts between graphenesheets, more conductive paths at a given filler concentration andtherefore higher conductivity.

It is difficult to compare the conductivity of composites with differentfillers and matrices as aforementioned factors can be different fordifferent systems. Even for graphene-polymer nanocomposites, thedifference in the dielectric and thermal properties of the matrix canlead to different tunneling resistance. The morphology, as well as thedispersion of graphene can be different depending on the functionalgroups and defects on graphene.

Preferential orientation of FGS during the shear molding process led tofewer contacts in the transverse direction and more contacts in thelongitudinal direction, causing the observed anisotropy in thenanocomposite conductivity.⁶³ ⁶³ Du, F. M.; Fischer, J. E.; Winey, K. I.Physical Review B 2005, 72, (12), 4

3.3 Mechanical Properties of Graphene-SE Nanocomposite.

The stress-strain curves of unfilled and FGS-filled SE are shown inError! Reference source not found., and the values of modulus, tensilestrength and elongation at break at various FGS weight and volumepercentages are shown in Table 1.

TABLE 1 Values of the modulus, tensile strength and elongation at breakfor FGS-SE nanocomposite at various FGS loadings. To convert from weightpercentage to volume percentage, SE density of 0.97 g/cm³ and FGSdensity of 2.25 g/cm³ were used. FGS FGS Modulus, E Elongation atStrength, ε_(b) weight % volume % (MPa) break, ε_(b) (%) (MPa) 0 0 1.33± 0.12  74 ± 16 0.57 ± 0.09 0.05 0.022 1.42 ± 0.13  66 ± 11 0.52 ± 0.060.2 0.088 1.64 ± 0.17 138 ± 17 0.90 ± 0.09 0.5 0.22 1.77 ± 0.11 139 ± 151.30 ± 0.10 0.8 0.35 1.93 ± 0.12 149 ± 40   1.49 ± 0.0.36 1 0.43 2.13 ±0.13 138 ± 19 1.78 ± 0.21 3 1.31 4.86 ± 0.44 112 ± 9  3.43 ± 0.22

For each FGS loading, the crosslinker and the catalyst concentrationwere chosen to yield the highest tensile strength of the samples. At0.05 wt. % (0.022 vol. %) FGS loading, no improvement in mechanicalproperties was observed. At 0.2 wt. % (0.086 vol. %), a 23% increase inthe modulus, an 87% increase in the elongation at break and a 58%increase in the tensile strength were observed. At 0.5 wt. % (0.22 vol.%) FGS, a 33% increase in the modulus, 87% increase in elongation atbreak and a 128% increase in the tensile strength were achieved. At a 3wt. % (1.34 vol. %) FGS loading, the modulus increased by 265%, theelongation at break increased by 51% and the tensile strength increasedby over 500%. Above the percolation threshold, the modulus and tensilestrength increased with the FGS concentration whereas the elongation atbreak increased initially with FGS content up to 1 wt. % (0.45 vol. %)and then decreased at higher FGS loadings. There appeared to exist acritical FGS concentration between 0.5 wt. % and 1 wt. % beyond whichthe elongation at break of the composite started decreasing.

To understand the effect of FGS on the tensile properties of SE, moviesof the tearing process of unfilled and FGS-filled SE in an SEM wererecorded to reveal the failure mechanisms. Snapshots of the movies areshown in Error! Reference source not found.A-4F. As shown in FIG. 4A,when the unfilled SE was deformed, the notch gradually opened up. Untila certain level of stress was reached, tearing was initiated from thetip of the notch due to stress concentration and it immediatelypropagated across the specimen with little resistance, leading to thefailure of the specimen. Due to its lack of ability to crystallize understrain, SE does not possess mechanism to arrest or deflect tearing andtransfer the mechanical load to other parts of the matrix that is notsampled mechanically. When percolated FGS network was introduced, thefailure mechanism of SE was altered dramatically. At the initial stageof the deformation of 0.5 wt. % FGS-SE, the notch opened up (FIG. 4B).At a certain stress level, tearing was initiated from the tip of thenotch. However, unlike the case of the unfilled SE, tear propagation wasresisted and the sample did not fail upon tear initiation (FIG. 4C). Thepercolated FGS network introduced resistance for tear propagation. Uponfurther deformation, tear was further opened, followed by thecatastrophic failure of the sample (FIG. 4D). The observations aboveclearly illustrated the enhanced tear resistance in SE introduced by thepercolated FGS network.

SEM was also used to characterize the tensile-fractured surface ofunfilled and FGS-filled. For the unfilled SE (FIG. 4E), a few ridgeswere observed and majority of the torn surface was smooth, indicatingthat once tearing was initiated, it propagated across the entire samplewith little resistance or distortion. The fractured surface of 0.5 wt. %FGS-SE is shown in FIG. 4F. The bright spots with submicron lengthscales were the FGS. The morphology of the failure surface was quitedifferent from that of the unfilled SE. The torn surface had more ridgesthan the unfilled SE, indicating the distortion of tear propagation bythe presence of percolated FGS.

To quantify the degree of tearing in the unfilled and FGS-filled SE,mechanical hysteresis measurement was undertaken. Hysteresis loss infilled rubber has been attributed to covalent bond rupturing in thematrix,⁶⁴ viscoelasticity⁶⁵ and the breakdown of the filler networkstructure⁶⁶. While viscoelasticity induced hysteresis loss isrecoverable, covalent bond rupturing and the breakdown of filler networkstructure can lead to irrecoverable hysteresis loss. Given the nature ofFGS-FGS interaction to be weak van der Waals' force, the contribution ofFGS network breakdown to irrecoverable hysteresis loss in FGS-SE islikely to be small. Therefore, measurements of irrecoverable hysteresisloss provide a method to quantify the degree of tearing in unfilled andFGS-filled SE. ⁶⁴ Suzuki, N.; Ito, M.; Yatsuyanagi, F. Polymer 2005, 46,(1), 193-201⁶⁵ Roland, C. M. Rubber Chemistry and Technology 1989, 62,(5), 880-895⁶⁶ Yamaguchi, K.; Busfield, J. J. C.; Thomas, A. G. Journalof Polymer Science Part B-Polymer Physics 2003, 41, (17), 2079-2089

The mechanical hysteresis data is shown in Error! Reference source notfound. Unfilled SE showed little irrecoverable hysteresis loss afterrecovery as evidenced by the overlapping of the stress-strain curves offirst stretch and the stretch after recovery (Error! Reference sourcenot found.), suggesting few covalent bonds rupturing in the matrix. The0.5 wt. % FGS-SE sample showed an observable difference in the twostress-strain curves and a higher irrecoverable hysteresis loss comparedto the unfilled SE, as shown in Error! Reference source not found. Theirrecoverable hysteresis loss ratio is determined to be 5.5%. The aboveresults suggest a higher degree of covalent bond rupturing in the matrixcaused by the presence of FGS, confirming the introduction ofdistributed deformation by adding FGS.

Based on the above evidence, the reinforcement mechanism of SE by FGScan be envisioned. To improve the failure properties of SE, a mechanismto arrest or distort the tearing initiated from the intrinsic defects isnecessary. The arresting or distortion of tearing could only be achievedwhen the length scale of the filler was much larger than the lengthscale of the initial tear size, which is determined by the size ofintrinsic defects. An individual FGS has a lateral size of hundreds ofnanometers, much smaller than the intrinsic defect size inelastomers.^(44,45) Therefore, only percolated FGS network leads to asimultaneous increase in tensile strength and elongation at break of SE.During deformation of the nanocomposite, tearing is initiated from theintrinsic flaws in the matrix and arrested or deflected by the presenceof percolated FGS. The arresting or deflection leads to loadtransferring to the FGS and other parts of the unstrained matrix,leading to the enhancement of tensile strength. Deformation and tearingare also distributed across a larger portion of the matrix compared tounfilled SE and the opening of tears causes dilation within the matrix,leading to the observed elongation at break increase.

The necessity of a percolated FGS network to improve mechanicalproperties shown here is in sharp contrast to previous studies ofmultiwall carbon nanotube (MWNT) and carbon black (CB) filled SBR inwhich the increase in strength and elongation at break occurred prior toelectrical percolation.^(67,68) The MWNT was shown to have length up to5 μm, whereas CB can form agglomerates up to hundreds of microns.⁶⁹Agglomeration of those fillers can readily achieve a length-scale thatis larger than the critical flaw size and improves the tensileproperties. ⁶⁷ Bokobza, L. Polymer 2007, 48, (17), 4907-4920⁶⁸ Reffaee,A. S. A.; El Nashar, D. E.; Abd-El-Messieh, S. L.; Nour, K.Polymer-Plastics Technology and Engineering 2007, 46, (6), 591-603⁶⁹Kohjiya, S.; Kato, A.; Ikeda, Y. Progress in Polymer Science 2008, 33,(10), 979-997

3.4 Modeling of Mechanical Reinforcement in Graphene-SE Nanocomposite

To corroborate with the reinforcement mechanism demonstrated above andmore importantly, to understand the reversal in elongation at break, alattice-based model is used to study the deformation mechanics of FGS-SEnanocomposite.⁴³ The deformed lattice with or without FGS is shown inError! Reference source not found.A-6D. Individual FGS or percolated FGSwas represented by the black lines. The torn matrix was represented bygreen regions. In the unfilled SE, tearing was initiated from thedefects and propagated across the sample without much resistance due tothe lack of tear arresting mechanism (FIG. 6A). As shown in FIG. 6B,when individual unpercolated FGSs were present outside of the defects,tearing was initiated from the intrinsic defects and propagated acrossthe sample undeterred as in the case of unfilled SE. No interactionbetween FGS and tears was observed and failure properties of SE weretherefore not improved. When FGS was percolated, the percolated networkhad a length scale much larger than the intrinsic defect size. When 1.6vol. % percolated FGS was added, tearing was initiated and distorted orarrested by the presence of percolated FGS (FIG. 6C). Through arrestingor distorting of tearing, deformation was distributed to the strongerparts of the matrix as evidenced by the distributed tearing. When FGSconcentration was increased to 5.2 vol. %, a higher degree ofdistributed deformation can be achieved as indicated by the increasedamount of torn matrix (FIG. 6D).

The simulated and normalized stress-strain curves of unfilled andFGS-filled SE are shown in Error! Reference source not found.A-7C. Thetermination of the stress-strain curve indicated the strain at the peakstress level. Peak stress was defined as the tensile strength of thesample and the strain at the peak strength was defined as the elongationat break. The simulation reproduced qualitatively the experimentalstress-strain curves of unfilled and FGS-filled FGS. The modulus andstrength increased with FGS concentration whereas the elongation atbreak increased initially and decreased above 5.2 vol. %.

Analysis of the mechanical load carried by FGS demonstrated theexcellent load carrying capacity of percolated FGS.⁴³ Through arrestingor distorting of tearing, mechanical load was transferred to the FGS andstronger parts of the matrix, leading to a significant increase in thetensile strength of the composite. The most interesting observation wasthe increase in the elongation at break of SE. To gain more insight intothe mechanism responsible for the elongation at break increase, westudied the fraction of matrix torn and the strain of tears as afunction of FGS concentration. The fraction of matrix torn increasedwith FGS concentration, suggesting an increasing degree of distributeddeformation. The strain of tears decreased with FGS concentration, dueto the close proximity of percolated FGS suppressing the opening oftears. The increase in elongation at break with FGS concentration can beexplained by the dilation effect of tearing. As tear opened up duringstraining of the sample, the sample could be elongated more. Our modelshowed that the elongation at break of the nanocomposite was dominatedby two factors: the fraction of matrix torn and the strain of tears.When the FGS concentration was increased, the two factors were competingwith each other and the reversal effect on the elongation at break withincreasing FGS concentration was a result of the domination ofdecreasing strain of tears over increasing fraction of matrix torn.

The elongation at break increase can also be influenced by other factorssuch as the strain and deformability of the untorn matrix. However,those effects cannot be investigated due to the technical limitation ofthe model.

To demonstrate the superiority of FGS at improving the mechanicalproperties of SE, modulus, tensile strength and their relativeimprovement of all filled PDMS-based SE were plotted against the fillervolume fraction for the concentration range studied in the present study(Error! Reference source not found.A-8B and Error! Reference source notfound.A-9B). Only studies that used pristine SE (containing no fillers)as the base polymer were chosen for the analysis. For each type offiller, the sets of data with the largest improvement in mechanicalproperties were chosen for comparison.

To compare the modulus enhancement brought by FGS to that by otherfillers, the modulus and improvement in modulus (calculated by dividingthe difference in modulus between the filled and unfilled samples by themodulus of the unfilled sample) were plotted against filler volumepercentage in Error! Reference source not found.A-8B. In terms of therelative improvement of the modulus, FGS is comparable to or better thanall the reported fillers except for MWNT. One thing needs to be noted isthat the unfilled SE in that MWNT-SE study had a modulus of 0.14 MPa,almost an order of magnitude lower than that of the unfilled SE used inthe present invention.¹⁷ The superiority of FGS at enhancing modulus canbe attributed to the high aspect ratio plate-like geometry and the highsurface area (higher than all the fillers reported in previous studies)which enabled the low percolation threshold and offered extensiveinterfacial interactions with the matrix and a higher degree of loadtransferring.

The strength and the improvement in strength of the FGS-SE nanocompositewere compared with those of other filled-SE as a function of fillervolume percentage, shown in Error! Reference source not found.A-9B. Interms of the relative improvement in strength, FGS performed comparablyor better than all other fillers except for MWNT. In the case of fumedsilica and precipitated silica, which yielded similar strengthimprovement as FGS, the unfilled SE in that study had a tensile strengthof 0.075 MPa,¹⁵ much lower than that of the unfilled SE in the presentstudy (0.57 MPa). FGS-SE also has the highest tensile strengths in theconcentration range studied. The superior ability of FGS to strengthenthe matrix is believed to be attributed to: 1. its plate-like geometry,high aspect ratio and surface area, which provide a low percolationthreshold and large load transferring; 2. The distributed deformationintroduced by the percolated FGS allowed more regions of the SE matrixto carry loads.

3.4 Barrier Properties of FGS-SE Nanocomposite.

The multi-functionality of FGS as a filler lies in its ability tosimultaneously improve the mechanical and electrical as well as thebarrier properties of SE.

Oxygen and nitrogen permeabilities of unfilled and FGS-filled SE weremeasured and the results are shown in Table 2.

TABLE 2 Oxygen and nitrogen permeability of unfilled and FGS-filled SE.Permeability was reduced by half with 3 wt. % (1.31 vol. %) FGS.Permeability (Barrier) Sample O₂ N₂ Unfilled SE 555 266 1 wt. % (0.43vol. %) FGS 514 249 3 wt. % (1.31 vol. %) FGS 283 137

With the incorporation of 1 wt. % (0.43 vol. %) FGS, permeability forboth gases was reduced by 7%. When 3 wt. % (1.31 vol. %) FGS was added,permeability was reduced by half. The improvement was better than thatin the clay-filled SE.¹⁸

The reduction in gas permeability is believed to be attributed to thepresence of FGS acting as impermeable barrier and increasing thediffusion path for the gas.⁷⁰ Additionally, it has been suggested thatdue to the large interfacial area in the nanocomposites, the propertiesof the matrix, such as the fraction free volume, can be reduced andfurther decrease in the permeability can be achieved.⁷¹ PDMS can formhydrogen bonding with hydroxyl groups and the interaction between PDMSand FGS provides a modification of the matrix permeability and thereforethe overall barrier property of the nanocomposite. ⁷⁰ Nielsen, L. E.Journal of Macromolecular Science 1967, A1, (5), 929-942⁷¹ Wang, Z. F.;Wang, B.; N.; Zhang, H. F.; Zhang, L. Q. Polymer 2005, 46, (3), 719-724

Obviously, numerous modifications and variations of the presentinvention are possible in light of the above teachings. It is thereforeto be understood that within the scope of the appended claims, theinvention may be practiced otherwise than as specifically describedherein.

1. (canceled)
 2. A method for production of a nanocomposite compositioncomprising a silicone elastomer matrix and functionalized graphenesheets having a surface area of from 300 m²/g to 2,630 m²/g, comprising:dispersing functional graphene sheets (FGS) in a polar solvent to forman FGS suspension; combining the FGS suspension with a vinyl terminatedpolysiloxane; removing the polar solvent; combining the resultingmixture with a crosslinker and a hydrosilylation catalyst; curing theresulting mixture to provide the nanocomposite; wherein the functionalgraphene sheets have a loading of greater than 0.05 wt % based on totalnanocomposite weight; wherein the hydrosilylation catalyst is present ata concentration of about 367 ppm to about 5600 ppm; wherein thenanocomposite comprises a silicon hydride to vinyl molar ratio of about1.5 to about 2.1; and wherein the functional graphene sheets are presentwithin the nanocomposite in a continuous three-dimensional connectednetwork in a manner wherein individual functional graphene sheets havenanometer scale separation at contact point between individualfunctional graphene sheets.
 3. The method of claim 2, wherein the curingis performed at elevated temperature for a period of time from 1 to 48hours.
 4. The method of claim 3, wherein the curing temperature is about100° C.
 5. The method of claim 3, wherein the curing is performed for aperiod of time from 5 to 30 hours, 20 to 25 hours, or approximately 24hours.
 6. The method of claim 5, wherein the curing is performed for aperiod of time of approximately 24 hours.
 7. The method of claim 2,wherein the functional graphene sheets have a loading of from 0.5 to 3wt %, based on total nanocomposite weight.
 8. The method of claim 2,wherein the silane cross-linker is a member selected from the groupconsisting of tetrakis(dialkylsiloxy)silanes and poly(hydromethylsiloxane) crosslinkers.
 9. The method of claim 2, wherein the silanecross-linker is a tetrakis(dimethylsiloxy)silane.
 10. The method ofclaim 2, wherein the vinyl-terminated polysiloxane has a viscosity offrom 100 to 300,000 mPas.
 11. An article formed from a nanocompositeproduced by the method of claim
 2. 12. The article of claim 11, whereinthe article is formed by casting.
 13. The article of claim 11, whereinthe article is formed by molding.
 14. The article of claim 11, whereinthe article is a member selected from the group consisting of coatings,adhesives, sealants, flexible electrodes, actuators, pressure sensors,printed circuits, and electromagnetic interference shielding materials.15. A method for production of a nanocomposite composition comprising asilicone elastomer matrix and functionalized graphene sheets having asurface area of from 300 m²/g to 2,630 m²/g, comprising: dispersingfunctional graphene sheets (FGS) in a polar solvent to form an FGSsuspension; combining the FGS suspension with a vinyl terminatedpolysiloxane having a viscosity of from 100 to 300,000 mPas; removingthe polar solvent; combining the resulting mixture with a crosslinkerand a hydrosilylation catalysts, wherein the silane cross-linker is amember selected from the group consisting oftetrakis(dialkylsiloxy)silanes and poly(hydromethyl siloxane)crosslinkers, and wherein the hydrosilylation catalyst is a memberselected from the group consisting of chloroplatinic acid, elementaryplatinum, solid platinum supported on a carrier; platinum-vinylsiloxanecomplexes; platinum-phosphine complexes; platinum-phosphite complexes;Pt (acac)₂, wherein (acac) represents acetylacetonate group;platinum-hydrocarbon conjugates; platinum alcoholates; RhCl(PPh₃)₃;RhCl₃; Rh/Al₂O₃; RuCl₃; IrCl₃; FeCl₃; AlCl₃; PdCl₂.2H₂O; NiCl₂; andTiCl₄; and curing the resulting mixture to provide the nanocomposite,wherein the curing is performed at elevated temperature for a period oftime from 1 to 48 hours; wherein the functional graphene sheets have aloading of greater than 0.05 wt % based on total nanocomposite weight;wherein the hydrosilylation catalyst is present at a concentration ofabout 367 ppm to about 5600 ppm; wherein the nanocomposite comprises asilicon hydride to vinyl molar ratio of about 1.5 to about 2.1; andwherein the functional graphene sheets are present within thenanocomposite in a continuous three-dimensional connected network in amanner wherein individual functional graphene sheets have nanometerscale separation at contact points between individual functionalgraphene sheets.
 16. The method of claim 15, wherein the hydrosilylationcatalyst is present at a concentration of about 367 ppm.
 17. The methodof claim 15, wherein the hydrosilylation catalyst is present at aconcentration of about 1,000 ppm.
 18. The method of claim 15, whereinthe hydrosilylation catalyst is present at a concentration of about1,280 ppm.
 19. The method of claim 15, wherein the hydrosilylationcatalyst is present at a concentration of about 5,600 ppm.
 20. Themethod of claim 2, wherein the hydrosilylation catalyst is present at aconcentration of about 1,280 ppm.
 21. The method of claim 2, wherein thehydrosilylation catalyst is present at a concentration of about 5,600ppm.